U.S. patent application number 11/874361 was filed with the patent office on 2008-09-04 for processing method for the production of amorphous/nanoscale/near nanoscale steel sheet.
This patent application is currently assigned to THE NANOSTEEL COMPANY, INC.. Invention is credited to Daniel James Branagan, Michael Breitsameter, Joseph Buffa, David Paratore.
Application Number | 20080213517 11/874361 |
Document ID | / |
Family ID | 39314840 |
Filed Date | 2008-09-04 |
United States Patent
Application |
20080213517 |
Kind Code |
A1 |
Branagan; Daniel James ; et
al. |
September 4, 2008 |
PROCESSING METHOD FOR THE PRODUCTION OF AMORPHOUS/NANOSCALE/NEAR
NANOSCALE STEEL SHEET
Abstract
The present disclosure relates to an iron alloy sheet comprising
.alpha.-Fe, and/or .gamma.-Fe phases wherein the alloy has a
melting point in the range of 800 to 1500.degree. C., a critical
cooling rate of less than 10.sup.5 K/s and structural units in the
range of about 150 nm to 1000 nm.
Inventors: |
Branagan; Daniel James;
(Idaho Falls, ID) ; Buffa; Joseph; (Maitland,
FL) ; Breitsameter; Michael; (Providence, RI)
; Paratore; David; (Warren, RI) |
Correspondence
Address: |
GROSSMAN, TUCKER, PERREAULT & PFLEGER, PLLC
55 SOUTH COMMERICAL STREET
MANCHESTER
NH
03101
US
|
Assignee: |
THE NANOSTEEL COMPANY, INC.
Providence
RI
|
Family ID: |
39314840 |
Appl. No.: |
11/874361 |
Filed: |
October 18, 2007 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
60829988 |
Oct 18, 2006 |
|
|
|
Current U.S.
Class: |
428/34.1 ;
164/47; 428/219; 428/220; 428/332 |
Current CPC
Class: |
B22D 11/0622 20130101;
C22C 45/02 20130101; Y10T 428/26 20150115; Y10T 428/13 20150115;
C22C 33/003 20130101 |
Class at
Publication: |
428/34.1 ;
428/332; 428/219; 428/220; 164/47 |
International
Class: |
B32B 15/04 20060101
B32B015/04; B32B 1/08 20060101 B32B001/08; B22D 23/00 20060101
B22D023/00 |
Claims
1. An iron alloy sheet comprising: .alpha.-Fe, and/or .gamma.-Fe
phases wherein said alloy has a melting point in the range of 800
to 1500.degree. C., a critical cooling rate of less than 10.sup.5
K/s and structural units in the range of about 150 nm to 1000
nm.
2. The iron alloy sheet of claim 1 further including structural
units of one or more of the following: (a) about 5 to 100 Angstrom,
or (b) about 10 nm to 150 nm.
3. The iron alloy of claim 1 further including structural units of
equal to or greater than 1 micron.
4. The iron alloy sheet of claim 1 wherein said iron alloy further
comprises complex boride, complex carbide or complex borocarbide
phases.
5. The iron alloy sheet of claim 2 wherein said alloy comprises
about 50% by vol. or greater structural units in the range of about
5 Angstroms to 100 Angstroms.
6. The iron alloy sheet of claim 2 wherein said alloy comprises
about 50% by vol. or greater of structural units in the range of
about 10 nm to 150 nm.
7. The iron alloy sheet of claim 1 wherein said alloy comprises
about 50% by vol. or greater of structural units in the range of
about 150 nm to 1000 nm.
8. The iron alloy sheet of claim 1 wherein said iron alloy has a
hardness of about 100 kg/mm.sup.2 to 3000 kg/mm.sup.2.
9. The iron alloy sheet of claim 1 wherein said iron alloy has a
tensile strength in the range of 100,000 lb/in.sup.2 to 950,000
lb/in.sup.2.
10. The iron alloy sheet of claim 1 wherein said iron alloy has a
tensile elongation at room temperature in the range of 0.01 to
40%.
11. The iron alloy sheet of claim 1 wherein said iron alloy has a
tensile elongation in the range of 0.1 to 280% at temperatures
greater than room temperature.
12. The iron alloy sheet of claim 1, wherein said alloy may assume
the configuration of a bar, tube, or pipe.
13. The iron alloy sheet of claim 1 wherein said iron alloy has a
thickness in the range of about 0.1 mm to 30 mm.
14. A method of producing an iron alloy sheet comprising: melting
an iron alloy wherein said iron alloy has a melting point in the
range of 800 to 1500.degree. C., a critical cooling rate of less
than 10.sup.5 K/s; and cooling said iron alloy and forming a sheet
at a rate sufficient to produce structural units in the range of
about 150 to 1000 nm.
15. The method of claim 14 further including iron alloy having
structural units of one or more of the following: (a) about 5 to
100 Angstroms, or (b) about 10 nm to 150 nm.
16. The method of claim 14 further including iron alloy having
structural units of greater than or equal to 1 micron.
17. The method of claim 14 wherein said iron alloy component
further comprises phases selected from the group consisting of
complex carbide, complex boride, complex borocarbide and
combinations thereof.
18. The method of claim 14 wherein said method further comprises
the step of: devitrifying said iron alloy sheet.
19. The method of claim 14 wherein said cooling said iron alloy
comprises twin roll casting, strip casting or belt casting.
20. The method of claim 14 wherein said iron alloy component
comprises about 50% by vol. or greater structural units in the
range of about 150 nm to 1000 nm.
21. The method of claim 15 comprising about 50% by vol. or greater
of structural units in the range of about 10 nm to 150 nm.
22. The method of claim 15 wherein said iron alloy component
comprises about 50% or greater structural units in the range of
about 5 Angstroms to about 100 Angstroms.
23. The method of claim 14 wherein said iron alloy is undercooled
in the range of 500.degree. C. to 1000.degree. C.
Description
CROSS REFERENCE TO RELATED APPLICATIONS
[0001] This application claims priority to U.S. Provisional
Application No. 60/829,988 filed Oct. 18, 2006.
FIELD OF INVENTION
[0002] The present invention relates to a method for producing
amorphous, nanoscale or near nanoscale steel from glass forming
alloys, wherein the alloys may have an angstrom or near nano-scaled
microstructure. The alloys may be formed into sheet, plate or
strip.
BACKGROUND
[0003] Since Sir Henry Bessemer first patented the twin roll method
for the production of steel sheet directly from a liquid melt over
150 years ago, a number of alternate methods of steel production
have been developed. Until the 1950's, ingot slab production was
the standard practice where steel was poured into stationary molds
or casks. Starting in the late 1950's, conventional slab casting
through continuous casting was developed as a new route to improve
yield, quality, and productivity in the production of steel. It is
used to produce semifinished billet, bloom, or slab for subsequent
rolling in finishing mills. In 1989, another steel manufacturing
process was developed called thin slab casting which was first
implemented by Nucor Steel. The process has allowed the production
of steel slabs which are typically thinner than those produced by
continuous casting. In addition, the process has been cited as one
of the two most important developments of the 20.sup.th Century. In
1998, the twin roll strip casting process (i.e. Castrip.RTM.) was
developed by Nucor Steel. In the strip casting process, molten
steel is poured into a smooth sheet in one step at the desired
thickness without the need for subsequent and expensive rolling
operations. This is achieved by directing liquid steel through
nozzles which are aimed between the gaps of two 500 mm spinning
copper alloy casting rolls.
[0004] Conventional steel alloys solidify by what may be termed
conventional liquid solid transformation routes. By this route,
generally a small amount of liquid undercooling may be achieved
prior to nucleation, resulting in the formation of coarse
structure, due to rapid diffusion at elevated temperatures. Growth
of corresponding crystals occurs in a superheated liquid melt,
resulting in conventional growth modes such as dendritic or
cellular growth. While theoretically, any metallic element or alloy
may form a glass, conventional steels may not form glasses under
normal solidification conditions as the critical cooling rates for
metallic glass formation of conventional steels may be extremely
high and generally in the range of 10.sup.6 to 10.sup.9 K/s.
[0005] In such a manner, conventional steel processes are designed
to cover the challenges in solidification of existing steel alloys
but are not designed for the particular challenges and technical
hurdles found in solidifying glass forming steels. For example the
twin roll process may work well for conventional plain carbon
steel. This may be because the primary goal is to solidify the
material while the material passes through the rolls; maximizing
the total amount of heat removal may only be a minor or secondary
goal. Since conventional steel alloys may undergo cooling to a few
tens of degrees sufficient to solidify the melt, not much heat has
to be removed before the solidification occurs.
[0006] However, in glass forming systems, in order to avoid
crystallization, the undercooling may be from the melting point
down to room temperature. It should also be appreciated that a
sufficient level of undercooling may be from the melting point down
to the glass transition temperature (T.sub.g), since below the
fictive glass transition temperature diffusion may be so slow that
the effective kinetics allows almost a total cooling rate
independence. Thus, as discussed above, the total undercooling
necessary in conventional steels may generally be
.ltoreq.50.degree. C. but for glass forming steels, the total
undercooling may be much greater and may typically be in the
500.degree. C. to 1000.degree. C., range depending on the alloy
chemistry. Such undercooling has limited the maximum thickness of
the amorphous structures achievable. Particularly as the amorphous
structures solidify they may tend to have low thermal conductivity
hindering the removal of thermal energy from the interior of the
structure. Thus, solidification behavior in glass forming metallic
alloys may be significantly different than what is found in
conventional metal solidification.
SUMMARY
[0007] In exemplary embodiment, the present disclosure relates to
an iron alloy sheet wherein the alloy has a melting point in the
range of 800 to 1500.degree. C., a critical cooling rate of less
than 10.sup.5 K/s and structural units in the range of about 150 nm
to 1000 nm. The alloy may also include one or more structural units
in the range of about 5 to 100 Angstrom or about 10 nm to 150
nm.
BRIEF DESCRIPTION OF THE DRAWINGS
[0008] The above-mentioned and other features and advantages of
this invention, and the manner of attaining them, will become more
apparent and the invention will be better understood by reference
to the following description of embodiments of the invention taken
in conjunction with the accompanying drawings, wherein:
[0009] FIG. 1 illustrates a schematic diagram of an exemplary twin
roll casting process;
[0010] FIG. 2 illustrates a model continuous cooling transformation
(CCT) diagram showing the effect of the two stage cooling on
metallic glass formation for the twin roll casting process;
[0011] FIG. 3 illustrates a schematic diagram of an exemplary twin
roll casting rollers;
[0012] FIG. 4 illustrates a schematic diagram of an exemplary twin
belt casting process;
[0013] FIG. 5 illustrates a model CCT diagram showing the effects
of the two-stage cooling process as a function of solidifying a
liquid melt on a twin roll and twin belt caster; and
[0014] FIG. 6 illustrates a model CCT curve showing the effects of
twin belt casting length as a function total undercooling achieved
and its effect on the two stage cooling.
DETAILED DESCRIPTION
[0015] The present invention relates to a method of forming a near
nanostructure slab, strip, or sheet steel, out of iron based glass
forming alloys. Glass forming steel systems may be classified as
metallic/metalloid glasses, wherein relatively little to no
crystallization occurs within the metallic matrix. It should be
appreciated that in metallic/metalloid glasses associations of
structural units in the solid phase of the metallic/metalloid glass
may occur, i.e., the glass alloy may include local structural units
that may be randomly organized in the solid phase, wherein the
structural units may be in the range of 5-100 Angstroms. As the
local structural units become more organized, the structure units
may increase and may develop phases in the nanoscale, (i.e., 10-150
nm structures), and near-nanoscale regions, (i.e. 150-1000 nm
structures).
[0016] The alloy chemistries may include multicomponent
chemistries, such as chemistries that may be considered steels or
steel alloys. A steel alloy may be understood as an alloy wherein
the primary constituent (e.g. greater than 50% by weight) may be
iron. In addition to iron, an additional 3 to 30 elements may be
used as alloy additions. The alloy chemistry may include relatively
high concentrations of P-group elements, which are non-metallic and
may therefore not be able to form metallic bonds. They may
generally include a binary eutectic chemistry consisting of iron
plus boron, carbon, silicon, phosphorous and/or gallium. However, a
very high percentage of these elements may dissolve in the liquid
melt, in the solid glass and to a lesser percentage in the
crystalline phases. When dissolved, the P-group atoms may form
covalent bonds, tying up free electrons and act to fill
up/partially fill up the outer valence band. This may result in a
reduction of thermal conductivity, which may be comparable to the
range of thermal conductivity associated with ceramic materials,
i.e. between 0.1 to 300 W/m-K, including all increments and values
therein. Other alloy additions may include transition metals such
as chromium, molybdenum, tungsten, tantalum, vanadium, niobium,
manganese, nickel, copper, aluminum, and cobalt; and rare earth
elements including yttrium, scandium, and the lanthanides.
[0017] The melting points of the multi-component alloys may be
lower than those of conventional commercial steel alloys and may be
in the range of about 800.degree. C. to 1500.degree. C., including
all increments and values therein, such as 960.degree. C. to
1375.degree. C., 1100.degree. C., etc. In addition, the alloys may
be glass forming, which may have critical cooling rates for
metallic glass formation less than 10.sup.5 K/s, such as between
10.sup.0 K/s to 10.sup.4 K/s. The phases formed during
solidification may depend on alloy chemistry, processing conditions
and thermal history during processing. Exemplary alloys may contain
ductile phases like .alpha.-Fe and/or .gamma.-Fe along with complex
carbide, complex boride, and/or complex borocarbide phases based on
various stoichiometries such as M.sub.2(BC).sub.1,
M.sub.3(BC).sub.2, M.sub.23(BC).sub.6, M.sub.7(BC).sub.3 and/or
M.sub.1(BC).sub.1. M may represent any transition metal which may
be present within the alloy composition.
[0018] Nucleation of glass forming alloys may be inhibited by
allowing high undercooling prior to nucleation or the onset of a
phase transition. Undercooling may be understood as the lowering of
the temperature of a liquid beyond the freezing temperature and
still maintaining a liquid form. If the level of undercooling
obtained is below the fictive glass temperature, T.sub.g, then a
metallic glass structure may be achieved. The fictive temperature
may be understood as the thermodynamic temperature at which the
glass structure may be in equilibrium. Thus, the total undercooling
may be in the range of 500.degree. C. to 1000.degree. C. depending
on the alloy chemistry, including all ranges and values
therein.
[0019] Accordingly, nucleation inhibition may occur if the critical
cooling rate of metallic glass formation is lower than the average
cooling rate of the manufacturing process of the steel alloy. In
addition, where nucleation may be at least partially avoided or
inhibited, latent heat related to the initiation of nucleation may
be reduced or not released. Thus, temperature increases due to
nucleation may be minimized, avoiding devitrification and/or
avoiding inducing a two-phase liquid/solid region, which may then
allow for solidification under conventional nucleation and growth.
The metallic glass may exhibit microstructural refinement including
an angstrom scaled microstructure. The glass sheet may then be
transformed into a nanoscale composite microstructure by a post
processing devitrification heat treatment.
[0020] The glass forming alloys may be processed using
manufacturing approaches such as twin roll casting, strip casting,
belt casting, etc., resulting in the development of microstructure
scales much finer than conventional steel alloys. Note that the
microstructures may include associations of structural units in the
solid phase that may be randomly packed together forming an
amorphous phase. The level of refinement, or the size, of the
structural units may be in the angstrom scale range (i.e. 5 .ANG.
to 100 .ANG.) if a metallic glass is formed; if nucleation or
crystallization is initiated, the level of refinement may include
the nanoscale region (i.e. 10 to 150 nm) and just above the
nanoscale range, that is "near nanoscale," (i.e. 150 to 1000 nm).
It should therefore be appreciated that the alloy may result in a
component that may include structural units in the range of about 5
.ANG. to 100 .ANG., 10 nm to 150 nm or 150 nm to 1,000 nm, as well
as combinations thereof. Accordingly, structural units in the range
of about 5 .ANG. to 100 .ANG., 10 nm to 150 nm or 150 nm to 1,000
nm may all be present in the iron alloy component. Furthermore,
structural units in the range of about 5 .ANG. to 100 .ANG., 10 nm
to 150 nm or 150 nm to 1,000 nm, may be present almost exclusively,
i.e., at levels greater than 90% by vol.
[0021] It should be appreciated that the level of refinement or
microstructural scale of the structural units may be determined by
various forms of X-ray diffraction with Scherrer analysis to
analyze peak broadening, electron microscopy (either scanning
electron microscopy or transmission electron microscopy) or Kerr
Microscopy utilizing a confocal scanning microscope. For example,
scanning electron microscopy (SEM) may be used to produce an
electron backscattered diffraction image, by detecting
backscattered electrons which may detect the contrast between areas
with different chemical compositions. Such an image may be used to
determine the crystallographic structure of a specimen. In
addition, SEM electron diffraction may be utilized. While the
spatial resolution of an SEM may depend on the size of the beam,
the resolution may also be dependent on the interaction volume, or
the extent of material which may interact with the electron beam.
In such a manner, the resolution may be in the range of about 1 to
20 nm.
[0022] Transmission electron microscopy (TEM) may also be used to
measure the microstructural units using techniques such as selected
area diffraction, convergent beam diffraction and observation with
or without rocking the beam. As it may be difficult to see the
short range order/extended short range order arising from molecular
associations due to the extremely fine ordering in metallic
glasses, advanced TEM techniques may be used. Dark field
transmission electron microscopy may be utilized as well as high
resolution transmission electron microscopy or field emission
transmission electron microscopy. Furthermore, scanning
transmission electron microscope may be used with aberration
correction software to produce images on the sub-Angstrom
scale.
[0023] Magnetic techniques such as direct measurements of domains
using a confocal scanning Kerr microscope may be employed to
measure domain size as well. Further measurements may also include
indirect measurements of nearest neighbor associations leading to
magnetic moments, Curie temperature, and saturation
magnetization.
[0024] In addition, the iron alloy may include 50% or greater by
volume (vol.) structural units in the near-nanoscale or in the
range of about 150 nm to 1,000 nm, including all values and
increments therein. It may also include about 50% or more by vol.
of structural units in the range of about 5 .ANG. to 100 .ANG..
Furthermore, the iron alloy may include about 50% or more by vol.
of structural units in the range of about 10 nm to 150 nm.
Furthermore, the alloy may include structural units in the micron
size range, i.e., greater than or equal to about 1 micron.
[0025] The properties and/or combination of properties found in the
near nanoscale alloy and slab, strip, or sheet produced there from
may be outside the existing boundaries of conventional steel sheet
and may include extremely high hardness, extremely high tensile
strength, superior strength to weight ratios, and enhanced
corrosion resistance.
[0026] In an exemplary embodiment, glass forming steel alloys may
be processed by techniques wherein the alloy may rapidly solidify,
which may be understood as cooling the liquid steel in a short
period of time so as to retain a microstructural scale which is
reduced in size. For example, rapid solidification may be obtained
by processing liquid steel on a metal chill surface that may
include a high conductivity metal such as a copper, copper alloy,
silver, etc. As alluded to above, exemplary rapid solidification
techniques include but are not limited to twin roll casting, strip
casting, and belt casting, such as horizontal single belt casting.
Steel strip, slab, or sheet components may be produced at the
minimum number of processing steps and at the lowest practical
thicknesses as possible. In an exemplary embodiment, there may be
no subsequent rollering stages. Solidified sheet may be understood
herein as having, e.g., a thickness from about 0.1 mm to 30 mm in
thickness including all increments and values therein, such as 0.5
mm to 15 mm thick, 10 mm thick, etc. Accordingly, by way of
example, sheet steel herein may be understood as a sheet of steel
having a length and width and the indicated thickness values. Such
length and width values may be in the range of 1 to 100 inches wide
and 1 to 1000 inches long, including all values and increments
therein. In addition, components such as tubes, pipes, or bars may
be formed as well.
[0027] In an exemplary embodiment, horizontal single belt casting
may be utilized wherein a chill surface is provided such that the
alloys may remain in contact with the single chilled belt for a
desired duration, depending on the length of the belt and roll
speed. Accordingly, the bottom fraction of the sheet next to the
chill surface may form a glass and the top surface may cool much
slower as it cools via radiation and natural convection. Thus, the
surface removed from the belt may crystallize at a much lower
amount of undercooling, which may result in a release of latent
heat. The release of latent heat may then cause a dramatic
temperature rise (i.e. recalescence), crystallizing a portion of
the underlying liquid melt. It should be appreciated that the
increase in temperature may be sufficient to bring the alloys to
the liquid region causing localized melting. Accordingly, it may be
appreciated that the single chilled belt procedure may only provide
relatively reliable glass formation for the bottom fraction and a
gradient of differing morphology proceeding to the outer
surface.
[0028] In another exemplary embodiment, twin roll casting may be
utilized wherein the melt may cool rapidly on the rolls.
Illustrated in FIG. 1 is a schematic diagram of an exemplary
embodiment of a twin roll casting system and method 10. As shown,
the liquid steel melt alloy 12 may have a first relatively high
temperature prior to contacting the primary cooling rollers 14.
When in contact with the rollers, which may be for example copper
alloys rollers, the alloy may cool very fast (i.e. conductive) at a
first rate R.sub.1 and may leave the wheel at a second relatively
high temperature T.sub.2, which may be somewhat less than the first
relatively high temperature T.sub.1. After leaving the chill
surface, the rate of heat removal may be relatively less than that
exhibited at the chill surface (i.e. radiative or naturally
convective) and results in a reduced cooling rate R.sub.2. The melt
may thus be solidified into a strip or sheet 16 and may pass
through secondary rollers 18. Thus, the cooling rate in twin roll
casting may be characterized as a two stage process.
[0029] The effects of two stage cooling are shown on the model
continuous cooling transformation (CCT) diagram for metallic glass
forming steel alloys shown in FIG. 2, wherein the C-Curve D
represents is the glass to crystalline transformation region and E
represents the glassy region. As shown, the initial cooling curve C
is rapid and in the range of possible development of glass forming
steel chemistries. However, the total amount of heat removal may be
insufficient and the liquid melt may come off the wheel in a
moderately undercooled condition at A. The much slower cooling rate
B of the liquid melt once removed from the wheel may result in the
formation of relatively larger crystals (i.e. >10 .mu.m) since
the nose of the glass to crystalline transformation (point F) is
may almost be entirely avoided. In FIG. 2, it should be appreciated
that Ts refers to the superheat temperature, Tm refers to the
melting point of the alloy, Tu.sub.1 refers to undercooling
temperature 1 at point A, Tu.sub.2 refers to undercooling
temperature 2, and Tg refers to the glass transition
temperature.
[0030] FIG. 3 illustrates another exemplary embodiment of twin roll
casting process 10. The rolls 14 may be counter-rotating forming a
nip through which the liquid alloy 12 is passed. Upon passing
through the nip and by contact with the rolls the alloy begins to
solidify along the roll surface and is brought together to form a
solid strip 16. Also, as shown is the total effective chill surface
(represented in phantom by arc S), which may be less than or equal
to one fourth of the roll circumference. For example, for a 500 mm
diameter roll results in only 393 mm (15.5'') of total chill
surface for the roll. Accordingly, it should be appreciated that by
increasing the diameter of the chill roll, the roll may exhibit a
larger surface area. However, the total chill surface may still be
approximately one forth of the roll circumferences.
[0031] In another exemplary embodiment, a twin belt may be utilized
as shown in FIG. 4. In this approach, two chill surfaces may be
provided which may allow for cooling of the alloy from both sides.
The total chill surface 20 (encompassing both the surfaces of the
top and bottom rolls forming the nip) may be much larger, i.e.
longer, and varied in length. The twin belts may be made out of
high melting point steel or highly conductive metals such as
copper, silver, gold or alloys derived from these elements. The nip
portion or entirety of the twin belts may be cooled using water or
other suitable coolant. The belts may be arranged in a horizontal
fashion (at an angle of 0.degree.) as shown or at an angle up to
vertical, such an angle in the range of +/-1 to 180.degree.,
including all increments and values therein. In addition, the belts
may be adjusted so as to provide constant pressure on the alloy as
it cools through out the forming processes, as the cooling alloy
may tend to shrink. In such a manner, the distance D (illustrated
by the phantom line) between the belt surfaces may be reduced along
the belt length L.
[0032] As illustrated in FIG. 5, the liquid melt may undergo single
stage cooling if the melt remains on the chill surface of the belts
for a sufficient period of time, such that the initial cooling
represented by curve C is rapid and the cooling rate is high. The
total length of the belts may be adjusted so that the liquid melt
comes off at a temperature where metallic glass precursors may be
formed. If metallic glass precursor sheet is formed, it can then be
transformed through various relaxation, recovery, single stage, and
multiple stage heat treatments into specific nanoscale structures
with a range of targeted sets of properties. Ideally, and as
illustrated at G, the point of melt removal would be at the glass
transition temperature Tg so that the second stage slow cooling
would not cause nucleation.
[0033] As illustrated in FIG. 6, the longer the chill belt, the
longer the liquid melt may undergo rapid cooling represented by
curve C. As the total belt length is increased, more heat can be
removed allowing for an ever greater of undercooling before the
sheet is removed. Achieving a much higher level of undercooling
would then better enable for the production of amorphous sheet,
plate, or strip. Accordingly, the longer the belt the less
secondary cooling may occur, represented by lines B, G, H, and I
wherein B represents the secondary cooling for a belt of a first
length L.sub.1, G represents a belt of a second length L.sub.2, H
represents a belt of a third length L.sub.3 and I represents a belt
of a fourth length L.sub.4, wherein
L.sub.1<L.sub.2<L.sub.3<L.sub.4. Note that even if the two
stage cooling does not avoid the nose of the CCT curve, such that
the cooling curve passes through the crystalline transformation
region, the higher undercooling would still allow the production of
nanoscale (i.e. 10 to 150 nm), or near nanoscale (i.e. 150 to 1000
nm) steel sheet, plate, strip, or other geometry.
[0034] Accordingly, the chill surface may be at a temperature that
is sufficiently low enough and exhibit a rate of heat flow that is
sufficiently high enough to prevent nucleation from occurring at
the surface and, preferably, throughout the thickness of the alloy.
In addition, it should be appreciated that while some nucleation
may occur, the microstructure size or growth may be limited to nano
or near nano scale.
[0035] Accordingly, if the critical cooling rate of the steel alloy
is higher than that of a given cooling process, the ability to form
a completely amorphous alloy may be compromised. However, due to
the glass forming nature of the alloys herein, high undercooling
may still occur prior to nucleation. Since nucleation may occur in
the glass forming alloys herein at several hundred degrees lower
undercooling than a conventional steel alloy, much greater
microstructural refinement may still occur. That is, although not
completely amorphous, relatively smaller crystalline domains may
still be formed with advantageous properties in those situations
where the critical cooling rate of the glass forming steel alloys
is higher than that of an applied cooling protocol. A lath
eutectoid may form in this case is one made up of alternating
platelets/laths with thickness's from 200 to 800 nm in size,
including all values and increments therein. A lath eutectoid may
be understood as alternating near nanoscale laths of ductile iron
and complex carbide phases such as borocarbide.
[0036] The properties produced from the steel may depend on a
number of factors including the level of microstructural
refinement, the microstructure that is produced and its constituent
phases, the glass forming steel alloy chemistry, the manufacturing
process chosen, the level of supersaturation, the post processing
conditions (if needed), etc. The contemplated macrohardness may be
approximately in the range of Rockwell C from 64 to 80, including
all values and increments therein. This hardness may be understood
to represent the hardness of the bulk which is an average of the
matrix and individual phases. The microhardness may vary depending
on the type of phases which are formed and may be approximately in
the range of HV 300 from about 100 kg/mm.sup.2 to 3000 kg/mm.sup.2
including all values and increments therein, such as 230 to 2500
kg/mm.sup.2, 850 to 2,000 kg/mm.sup.2. The contemplated tensile
strength may be in the approximate range of 100,000 lb/in.sup.2 to
950,000 lb/in.sup.2, including all values and increments therein
such as 170,000 lb/in.sup.2 to 480,000 lb/in.sup.2. The
contemplated tensile elongation at room temperature may be in the
approximate range of 0.01 to 40% including all values and
increments therein, such as 1 to 20%. At elevated temperatures,
such as those greater than room temperature, the contemplated
tensile elongation may be approximately in the range of 0.1 to 280%
including all values and increments therein, such as 4 to 60%.
Thus, the tensile elongation may be high at elevated temperatures
and may allow thermomechanical transformation (if necessary) of the
slab, strip, or sheet products into industrially usable shapes and
sizes.
[0037] The near nanostructured steel alloys may be used in a number
of applications. In one exemplary embodiment, the steel alloys may
be used in applications where there may be exposure to highly
corrosive or abrasive environments. The alloys may therefore be
used to replace or in combination with nickel base superalloys,
(i.e. 625, C-22) or stainless steels (i.e. 316, 304, 430, etc.).
The steel may be used as or may assume the configuration of a wear
plate which may be used as a replacement for or in combination with
conventional high hardness sheet material like tool steel, Hardox,
Brinell 500, etc, or weld overlay wear plates such as those
hardfaced with chrome carbide, WC, complex carbide, tungsten
carbide etc. The wear plate produced may have wide applicability in
the heavy construction, mining, and material handling industries in
a number of applications including but not limited to chutes,
ground engaging tools, truck beds, undercarriage components etc.
Additional uses of the near nanostructured sheet may include
aerospace applications, steel armor or military armor plate,
protecting infrastructure, civilian vehicles and military vehicles,
wherein the alloys may be used to replace or in combination with
titanium alloys, ultra high strength steel, ceramic materials,
conventional armor steel or reactive armor steel etc.
[0038] The foregoing description is provided to illustrate and
explain the present invention. However, the description hereinabove
should not be considered to limit the scope of the invention set
forth in the claims appended here to.
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